Thermomechanical strengthening of the superalloys

ABSTRACT

The superalloys are strengthened in a process involving both thermal and deformational treatments under controlled conditions. The method is particularly effective for the nickel-base superalloys of the gamma - gamma &#39;&#39; -type having a volume fraction of the gamma &#39;&#39; phase exceeding about 25 percent at room temperature, and for the superalloys precipitating the topologically close-packed phases such as the sigma phase. It relies on the establishment of a microstructure wherein the gamma &#39;&#39; phase is precipitated in a uniformly distributed array having an interparticle spacing not exceeding about 5 microns; warm working the alloy to effect an area reduction of at least 15 percent; and subjecting the alloy to a stabilization heat treatment. The strength increase is attributable to a particular thermally and mechanically stable array of microcrystalline imperfections thus established and, in those alloys precipitating the sigma phase, also by an altered sigma phase morphology.

United States Patent Owczarski et al.

[54] THERMOMECHANICAL STRENGTHENING OF THE SUPERALLOYS [72] inventors: William A. owczarski, Cheshire; John M.

Oblak, Rocky Hill, both of Conn.

[ 51 Feb. 15, 1972 Primary Examiner-Richard 0. Dean Attorney-Richard N. James ABSTRACT The superalloys are strengthened in a process involving both thermal and defonnational treatments under controlled conditions. The method is particularly effective for the nickel-base superalloys of the 'yy'-type having a volume fraction of the 7' phase exceeding about 25 percent at room temperature, and for the superalloys precipitating the topologically closepacked phases such as the sigma phase. it relies on the establishment of a microstructure wherein the 7 phase is precipitated in a uniformly distributed array having an interparticle spacing not exceeding about 5 microns; warm working the alloy to effect an area reduction of at least 15 percent; and subjecting the alloy to a stabilization heat treatment. The strength increase is attributable to a particular thermally and mechanically stable array of microcrystalline imperfections thus established and, in those alloys precipitating the sigma phase, also by an altered sigma phase morphology.

16 Claims, 7 Drawing Figures strengthening,

THERMOMECHANICAL STRENGTHENING OF THE SUPERALLOYS This application is a continuation-in-part of application Ser. No. 746,013 filed July 19, 1968 and now abandoned.

BACKGROUND OF THE INVENTION The present invention is most conveniently characterized as a metal processing technique and it is particularly adapted to improving the mechanical properties of the nickel-base and cobalt-base superalloys.

The superalloys are, in general, those alloys which display very high strengths at very high temperatures and, thus, which have significant utility in the fabrication of gas turbine engine components. The typical nickel-base superalloy of this type, for example, is essentially a nickel-chromium solid solution (7 phase) hardened by the additions of elements such as aluminum and titanium to precipitate an intermetallic compound (7' phase). The usual intermetallic compound, which is represented by the formula Ni (Al, Ti), is an ordered facecentered-cubic structure with aluminum and titanium at the corners of the unit cell and nickel at the face centers. These alloys also normally contain cobalt to raise the solvus temperature of the 7' phase, refractory metal additions for solution and carbon, boron and zirconium to promote ductility and fabricability. In the monocrystal form, these alloys may have reduced quantities of carbon to prevent crackproneness associated with the formulation of MC-type carbides.

One series of alloys to which the present invention has particular applicability are those nickel-base superalloys having a quantity of the 'y' precipitate exceeding about 25 volume percent at the hot working temperatures and which is stable within the matrix at this same temperature. Representative of the alloys of this type are those identified in the industry as follows:

Designation Composition (by weight) MAR-M200 Conventional alloys of the type described above find extensive usage in the gas turbine engine industry. As part of the continuing development programs, improvements in the properties of the superalloys are sought with or without significant alterations of the alloy chemistry. One very common problem with the typical superalloy, which usually represents a balance between strength and oxidation resistance among other criteria, relates to its inherent susceptibility to damage by such phenomena as high cycle and low cycle fatigue. Accordingly, a number of the development programs are directed toward the improvement of such specific properties without diminution of the other advantageous physical characteristics of the alloys in question. The thermomechanical strengthening sequence of the present invention is the outcome of one such program.

In a copending application of the same assignee entitled Fabrication Method for the High Temperature Alloys, Ser. No. 692,705, filed Dec. 22, 1967, now US. Pat. No. 3,519,503 a simultaneously-applied elevated temperaturecompressive working operation was described to impart forgeability to the advanced gas turbine alloys. Similarly, in an other copending application of the same assignee entitled Method to impart Fabricability to the Nickel-Base Superalloys, Ser. No. 745,958, filed July 19, 1968, now abandoned the same problem with the superalloys, namely a lack of fabricability, was attacked by establishment of a particular polygonal subcell microstructure as attained by heat treatment and working. The present invention utilizes somewhat similar processing, particularly in the early sequences of the method, but ultimately achieves a thermally and mechanically stable array of microcrystalline imperfections for strengthening purposes.

Another problem often encountered with the superalloys is the formation of undesirable phases during exposure to elevated temperatures or to some particular operating regime. The phase known as the sigma phase (0') is a particularly well known example of a precipitate of this general type and many of the commercially available superalloys, such as [NCO 901 and [NCO 718, are specifically formulated to avoid the formation of these undesirable phases. Typically, the sigma forms as hard, brittle platelets which provide the natural sites for mechanical weakness and, in addition, compete with the 7 matrix phase for the solid solution strengthening elements. A detailed discussion of the sigma phase may be found in an article by E. 0. Hall et al. The Institute for Metals, Vol. 1 1 (1966) p. 61.

Sigma belongs to a class of intermetallic phases identified as topologically close-packed (TCP) phases which typically form in a platelike morphology in a size comparable to the alloy grain size. These TCP phases, such as the sigma, chi or mu phases, occur in both the nickel-base and cobalt-base alloys.

It has been shown that the sigma phase precipitate in solution strengthened stainless steels may be changed from the detrimental platelet morphology to a less harmful globular shape by cold working. See, for example The Effect of Sigma Phase on the Short-Time High Temperature Properties of 25 Chromium-20 Nickel Stainless Steel, Guarneiri et al., Trans. ASM, Vol. 42 (1950) p. 981. However, the use of cold work for this purpose in the nickel-base superalloy systems, for example, is generally not advisable. First, the cold-worked microstructure is inherently unstable in these superalloys at their representative service temperatures. Secondly, the distribution of deformation in the cold-worked alloys and, consequently, the sigma phase configuration is apt to be quite heterogeneous. Furthermore, in the superalloy field cold work is normally difficult to achieve following full aging and frequently introduces embrittlement or other microstructural instabilities if provided prior to aging.

SUMMARY OF THE INVENTION The present invention provides a method for improving the yield and tensile strengths and improving the creep and fatigue resistance of the advanced nickel-base superalloys by thermomechanical means. The improved properties are provided by a stable array of microcrystalline imperfections established by controlled heat treatment and deformation, and preserved and stabilized by subsequent heat treatment causing further 'y' precipitation and in the case of the alloys prone to sigma formation, a finely dispersed precipitate of equiaxed sigma.

Basically, the method disclosed herein involves the following procedure:

1. An initial microstructure is established by heat treatment involving solutioning plus aging to provide a microstructure consisting of a uniformly distributed y phase having an effective interparticle spacing not exceeding about 5 microns and a volume fraction not less than about 25 percent.

2. The material so heat treated is deformably processed at elevated temperatures subject to the following conditions:

a. the deformation temperature is selected to preserve the volume fraction and distribution previously established and usually corresponds to the aging temperature of step l b. the deformation temperature must not exceed either the solvus temperature of the y phase or the gross 3 ,642,543 3 4 recrystallization temperature and must not be less then D RIP I N F THE PREFERRED EMBODIMENTS the minimum recovery temperature;

c. total deformation must be equivalent to at least about a percent reduction (but is not likely to exceed a reduction of about 60 percent).

. Following deformation the material is heat treated at a temperature which must not exceed the temperature of deformation and prior aging. This final heat treatment at lower temperatures is utilized to stabilize the microdefect array and additionally strengthen the alloy by normal Throughout the description reference is made to various alloys, temperatures, heat treatments and other process parameters. As utilized, the terminology employed herein is generally consistent with that utilized in the art. The following table sets forth certain select properties for several of the advanced superalloys which are of particular interest in the gas turbine engine industry and to which the methods of the present inven- 10 tion are particularly applicable.

precipitation hardening. in this final heat treatment there TABLE I is usually precipitation of additional 7 and, as appropriate, precipitation of the TCP phases. For certain al- PROPERTIES O E SUPERALLOYS loys, a single heat treatment may suffice; for other alloys multiple heat treatments may be preferred, depending l5 Appmimhc th ti re i .tat. t t f th melting 7' recovery 7' aging 3 Op mum p C F ures 6 Alloy point solvus temperature temperature various phases and the properties desired in the finished product.

f h b d MAR-M200 2300 2250 1850 isso-zioo I A signl icant measure 0 strengt emng may e ac ieve m [M900 2300 2200 1850 18504050 either of two microcrystallme forms, the generation of which Udimet 700 2220 2100 1800 1800-1975 in the alloys of the y-y type is dependent upon the particular temperature within the aging temperature range at which the 'Coincidentwith the uppertemperaturclimit for planarslip aging and deformation is accomplished. The more significant strengthening is achieved by working the alloy nearer the lower end of the aging temperature range, i.e., nearer the Preferred, in context minimum recovery temperature. Processing at this tempera- I will e understood t e ab e e p es a ture results in a warm-worked substructure comprising a ranrep es nt tiv r n Of h above Parameters ing d pe domly nonoriented, homogenous dislocation distribution. dent to some extent on the e a t a oy comp ition a d p i r When the material is deformed at a temperature nearer the history.

upper end of the aging temperature range, i.e., nearer the 501- The representative conditions utilized to establish specific vus temperature ofthe y precipitate,apolygonal substructure substructures in certain of the superalloys are set forth in is achieved and the dislocations are mobile enough to align T l TABLE IIENOMINAL CONDITIONS TO ESTABLISH SPECIFIC SUPERALLOY SUBSIRUCTURE F-) Pro-work heat treatment Working Post-work temperature heat Alloy Substructure Solution Age range treatment Udimet 700 Polygonal 2,140(r)l2,160/ 1,925-1,975/ 1, 926-1, 976 Normal stabilize and final age.

rs. 4 hrs. Udimet 700 Warm worked 2,102-2160/ l,& l10?1,850/ 1, BOO-1,860 D0.

rs. rs. MAR-M200 Polygonal 2,240ll)2,250/ 1,97l}i]-2,025/ 1,975-2,026 Do.

trs. 4 rsv MAR-M200 Warm worked..." 2,240g-2,250/ 1,23%1350/ 1, 900-1, 950 D0.

rs. rs.

No'rE.Times are those utilized for convenience and represent no limitation on the maximum or minimum allowable times.

Aging and working at temperature levels between the bands listed above will likely produce mixtures of substructure types.

themselves at subcell boundaries to provide a regular array of In performing a solution heat treatment, the conditions are defects along these boundaries. normally selected to dissolve the maximum quantity of the y In some cases, by processing the material at one temperaprecipitate into solid solution. In the case of materials like ture condition followed by subsequent processing at another Udimet 700, all of the 7 phase can be solutioned. With condition, or by processing at intermediate temperatures, it is MAR-M200 most, but perhaps not all, of the 7' phase is dispossible to produce both kinds of dislocation arrays in the solved. Solution heat treatments are conducted near or above micfostfuctufe the solvus temperature for the particular alloy involved but below the solidus temperature or that at which hot short- BRIEF DESCRIPTION OF THE DRAWING ness occurs. in the case of the highly alloyed materials such FIG. 1 is a photomicrograph taken utilizing electron as MAR-M200 r where the solvus terhpel'mlh'e approaches microscopy techniques showing the microstructure of a Sohdhs temperature solhhonihg y actually be done Slightly Udimet 700 alloy sample subjected to a conventional below the solvus temperature- Strengthening h treatment M 24,0QOX) Aging of the alloy results in precipitation of the 7' phase. In

FIG 2 i a photomicrograph f a Udi 700 alloy Sample the context of the present invention, aging may be undertaken processed to produce the polygonal substructure of the at any temperature above the mihhhhm recover) temperature present invention (24,Q()() or that temperature above which nonplanar slip occurs, and

HO. 3 is a photomicrograph of the same alloy processed to below the SOh/hs temperature of the 7' p Aging and Workproduce the warm-worked substructure of the present invenihg hear the upper end of this range Promotes formation of 3 tion, (24,000X) polygonal substructure, while aging and working near the FIG. 4 is a photomicrograph f a sigma prone ll lower end ofthe range results in a warm-worked microcrystalprocessed conventionally. (500x) line array. A further qualification on the aging parameters is FIG. 5 is a photomicrograph f h ll f FIG 4 processed established, particularly in a temperature-time relationship, to according to the present invention. (500X) The sections precipitate a minimum q y of the 7' Phase of about 25 shown in FIGS. 4 and 5 have been stain etched to reveal only lum P rcent a an interparticle spacing not exceeding the sigma and arbideplmses. about .5 microns.

NUS. 0 anti I are graphs computing the hardness oi the D l'orlmlilon is ur 'mhly undertaken M numlr-ml rm treatments prone alloys as conventionally proc ssed and n lure at which the alloy has been aged. However. it may he rm processed according to the present invention. dertaken at any elevated temperature within the aging range provided the microstructure established in aging is essentially maintained. The total deformation must exceed that equivalent to about a 15 percent reduction of area, and the strength increase which finally results is usually achieved in the range of a 15-60 percent reduction. This is not to say that further reduction cannot be made without deleterious effect, but rather that the maximum advantages in physical properties will have been achieved at the point where a 60 percent reduction has been achieved.

Deformation is normally accomplished with a 5-10 percent reduction per pass with reheat between passes to reestablish the temperature. With greater degrees of deformation, particularly when imposed in a limited number of passes, the synergistic effects of working plus external heat are necessarily considered It is desirable to prevent precipitation simultaneously with the working process, hence, the desirability of working the alloy at the same temperature at which it was aged is established. And, as previously mentioned, the selection of the aging-working temperatures is additionally influenced by the particular form of strengthening desired, i.e., the polygonal or warm-worked microstructure or combination thereof.

The final postwork heat treatment comprises a normal stabilization and precipitation aging. The temperature in the postwork heat treatment sequence must not exceed the temperature of deformation and final aging. Its purpose is to promote the final arrangement of the microcrystal imperfections into a stable array in the matrix through the further precipitation of the 7' phase and, in alloys containing more than about 0.05 weight percent carbon, to precipitate intragranular carbides. Both of these subsequent precipitation events further stabilize the microdefect array and additionally strengthen by normal precipitation hardening.

A number of specific materials were tested under various conditions. The following examples are representative of the preferred embodiments described herein. Further specific details relative to the present process will be evident to those skilled in the art, not only from the processing details set forth, but also from the evaluation of the various specimens so processed.

EXAMPLE 1 Polycrystalline Udimet 700 was thermomechanically processed to produce strengthening through achievement of a polygonal microstructure in the following manner:

Solution anneal at 2,140 F./4 hrs., air cool; age l,925 F./4 hrs.; air cool; work at l,925 F. to approximately a 60'percent reduction in area followed by stress relief at l,925 F ./2 hrs.; final age at l,550 F./4 hrs. and 1,400 F./16 hrs.; and slow cool to room temperature.

The physical property improvements were verified by various test procedures. These are set forth for Samples 1-5 in the following summary.

TENSILE PROPERTIES Test 0.2% yield Ultimate Percent temp., strength, strength, (in 1'0 Percent F. K s.l. K s.l. along.) R of A.

80 202 265 11.7 11.4 80 210 267 9.6 10. 6 80 140 200 17 20 1, 000 189 244 6. 7 9. 6 1,200 182 226 18.3 41 1,200 124 180 16 20 1,400 142 166 29 49 Normal U-700. 1, 400 120 150 33 40 Creep Properties Thermomechanically treated samples tested at 100 k.s.i., 1,200" F., showed 01 percent creep in 72 hours. The normal creep for the untreated alloy is 0.1 percent creep in about 20 hours. Stress Rupture Properties Thermomechanically treated samples were tested at 125 k.s.i., 1,200 F. and exhibited failure in 790 hours. Comparable times forthe untreated alloy average about 200 hours.

EXAMPLE 2 A monocrystal specimen formed from low carbon Udimet 700 was thermomechanically treated to produce a polygonal substrugture as follows: 7

Solution anneal at 2,l40 F./64 hrs.; age at 1,925 F.;4 hrs.; work at l,950 F. to 41 percent area reduction; then slow cool to room temperature.

With this specimen the 0.2 percent yield strength was 152 k.s.i. For the untreated alloy the 0.2 percent yield strength is 121 k.s.i.

A similar treated specimen, tested in low cycle fatigue at l,400 F. and 1.6 percent total strain amplitude, had a life of 3.093 cycles. The comparable lifetime of an untreated monocrystal of the same composition is about 900 cycles, that of an untreated polycrystalline sample of the same composition about 200 cycles.

EXAMPLE 3 A monocrystal sample of low carbon MAR-M200 was thermomechanically treated to produce a polygonal substructure as follows:

Solution anneal at 2,250 F./ hrs.; age at 2,000 F./2 hrs.; work at 2,000 F. to a 42 percent reduction followed by a slow cool to room temperature.

The 0.2 percent yield strength of this sample was 173 k.s.i. which compares to a k.s.i. yield strength for the untreated alloy. in low cycle fatigue at room temperature with a 1.5 percent strain amplitude, life was 3.038 cycles as compared to a lifetime of about 700 cycles for the untreated monocrystal.

EXAMPLE 4 A Udimet 700 polycrystalline specimen was thermomechanically processed to produce a warm-worked substructure as follows:

Solution anneal at 2,140 F./4 hrs.; age at l,825 F./4 hrs.; work at l,800 F. to a 30 percent area reduction (specimen 1 and a 60 percent area reduction (specimen 2); then age as shown.

The physical property improvements verified by the foregoing data were achieved without change in alloy composition by the provision of a specific substructure by thermomechanical treatment of the alloy. The substructure size is of the order of 5 microns or less and was achieved because the original particle size and distribution of the 'y' precipitate, as established by heat treatment, provided a network of particulate barriers with a spacing of 5 microns or less. The strength increase established is maintained as long as the substructure is maintained, and, hence, is evident up to at least the minimum recovery temperature for the particular alloy involved.

Either of the two types of dislocation structures previously mentioned while providing strengthening in and of themselves also provide means for significant strengthening in another fashion. When they are produced by working the superalloys at temperatures above the TCP phase precipitation range they inhibit the precipitation of these phases in the platelike Widmanstatten form, modifying the shape of the sigma, for example, to the more innocuous globular form and decreasing the actual precipitate particle size as clearly evidenced in FIGS. 4 and 5. The homogeneity of deformation prevents extensive nucleation of the sigma phase in the typical platelike mode and the stability of the dislocation structure prevents extensive recovery at the sigma aging temperature.

An alloy was designed to be highly sigma prone. Material was vacuum induction melted and cast with electrodes which were subsequently consumable arc melted under vacuum into two 1 inch diameter bars. Chemical analyses of both bars were performed using atomic absorption (Cr, Ti, Mo, Co, Al) spectroscopy (B, Zr) and combustion (C) methods. The results are shown in Table [11. W7, 77,.

The bars were solution annealed at 2, 150 F. for 4 hours and fast air cooled. Bar A was then aged at l,975 F. for 4 hours to precipitate the y phase. It was then swaged to a 60 percent reduction in area at 1.975 F. using a reduction of 6 percent per pass with a minute reheat between passes. No difficulty was experienced during the working operation. Bar B was given an 1,875 F. age for 4 hours. Swaging at 1,875 F. resulted in severe cracking during the first swaging pass. The l,875 F. age was found to promote the precipitation of both 7' and the platelike sigma.

TABLE III.COMPOSITION BY WEIGHT PERCENT Cr Tl M0 Co Al B Zr C N1 BarA 16.6 4 6 5.2 16.5 6.0 .005 .00l .025 1381. Bar B 16.6 4 6 5.2 16.3 5.1 .005 .001 .025 Be].

solution anneal 2,150 F./4 hours 2. age for l,975 F./4 hours 3. swage 60%RA at 1,975 F.

4. age l,200-l,875 F.; 4 or 24 hours For the purposes of comparison, samples were prepared which had an identical thermal history but no deformation processingv Specimens of this type are hereinafter referred to as standards.

The hardness of the warm-worked material was significantly higher than that of conventional nickel-base superalloys. As shown in FIG. 6, hardness values in the R 50s were obtained after aging at temperatures below 1,700 F. For comparison, the hardness of fully heated treated Udimet 700 is W R 39. In fact, the hardness of the sigma containing material is nearly equivalent to the low R 60 values of high-speed tool steel. For this alloy, a maximum hardness of R, 57.5 was obtained after a 24 hour age at 1,400 F. The Vickess hardness curves of FIG. 7 indicate that the peak hardness after a 4 hour age is nearer l,500 F. The shift in peak hardness to l,400 F. with increasing aging time indicates that any softening mechanisms, such as recovery or coarsening, occurring at l,400 F., are more then compensated, presumably by continued precipitation of the sigma phase.

As in the case of hardness, the yield strength of the warmworked material in compression was significantly higher than that of the conventionally processed nickel-base superalloys, as shown in Table IV. At room temperature, the ductility of the hardest warm-worked specimen is above that of tool steel but the yield strength is 100,000 p.s.i. lower. However, at 1,000 F. the strength of the tool steel has dropped significantly while the warm-worked superalloy has lost very little strength. At this temperature the yield strengths of the two alloys are about equal but the ductility of the tool steel is inferior.

TABLE IV Percent Testing R, Compressive Max. comdeform. temg hardyield strength, pressive at max. Specimen ness 0.2% offset strength strength A-l R.T. 67. 5 265, 000 314, 000 4. 1 M-1. 355,000 410,000 2.2 A-2 243, 000 302, 000 3. 8 M2.. 253,000 312,000 2.3

NorE.A1, A-2: warm-worked specimens; M-1, M-2: fully heat treated M-50 tool steel.

the present invention has revealed that not only may the detrimental effects of sigma be alleviated in the superalloys, but also significant improvements may be provided therefrom. The presence of substantial quantities of dispersed, equiaxed sigma in the superalloys may be utilized to enhance strength and to provide high hardness without embrittlement. This suggests that these modified alloys may have utility as high temperature bearing alloys for, while high speed tool steels are rarely used in bearing applications above 600 F. and are limited to temperatures below l,000 F., the equiaxed sigma superalloys are metallurgically stable to temperatures well over 1,000 F.

The significance of the invention is in fact considerably broader than a strengthening or hardening improvement may indicate. Alloy formulations may now be made without concern over any detrimental effects incident to sigma phase precipitation. Furthermore, not only may the sigma phase be utilized in an advantageous manner in the conventional sense but also, since its chemistry is variable, by changing its chemical makeup, control over its precipitation temperature can be exercised as necessary to tailor the alloy to a particular operating or processing regime. Still further, the controlled precipitation of the sigma phase or other topologically closepacked phases can now be utilized to advantage in the cobaltbase and austenitic iron-base alloys. In practical terms, therefore, a broad new field of alloy chemistry has been opened to practical utilization.

One significant advantage of the development, in addition to the strengthening effect, is of fundamental interest to the gas turbine engine industry. Because the tendency for the formation of sigma is no longer a factor in alloy composition, high aluminum and titanium contents can be used resulting in very low alloy densities. High aluminum and titanium contents promote copious precipitation of 7 which enriches the 7 matrix in sigma forming elements. The density of the principal tested alloy, for example, was 0.282 pounds per cubic inch which is lower than nearly every other nickel-base superalloy except lN-10O which is also sigma prone.

Although this invention has been described with great particularity and with reference to specific materials, processing parameters and examples for the purposes of illustration, the invention in its broader aspects is not limited to the specific details described, for obvious modifications will occur to those skilled in the art.

What is claimed is: l. The method of processing the superalloys subjected to the precipitation of intermetallic compounds including a 7 phase and, optionally, a topologically close-packed phase which comprises the steps of:

heat treating the alloy, including the step of aging above the minimum recovery temperature ofthe alloy, to establish a microstructure having the y precipitate in a stable homogenous distribution having an effective interparticle spacing not exceeding about 5 microns and a volume fraction of the precipitate not less than about 25 percent at the working temperature; warm-working the alloy to effect a reduction in area of at least 15 percent while maintaining essentially the same 7' phase morphology established in the prior heat treatment, providing a microdefect array of regular geometry; and

heat treating the alloy at a temperature not exceeding the working temperature to precipitate the 7' phase remaining in solution and to precipitate any topologically closepacked phases as globular particles, stabilizing the microdefect array and providing additional strengthenmg.

2. The method according to claim 1 wherein:

the initial heat treatment includes aging at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation upon warm-working of the regular microdefect array comprising a polygonal subcell structure.

3. The method according to claim 1 wherein:

the initial heat treatment includes aging at a temperature sufficiently close to the minimum recovery temperature of the alloy to promote formation upon warm-working of the microdefect array comprising a warm-worked metallurgical substructure consisting of a randomly nonoriented homogenous dislocation distribution.

4. The method according to claim 2 wherein:

the aging in the initial heat treatment is also above the precipitation temperature for the topologically closepacked phases.

5. The method according to claim 3 wherein: Y

the aging in the initial heat treatment is also above the precipitation temperature for the topologically closepacked phases. 6. The method of processing the nickel-base superalloys of the 7-7 type having a quantity of the 7' phase at room temperature exceeding about 25 volume percent and, optionally, a sigma phase precipitate which comprises the steps of: heat treating the alloy to solution the precipitate; aging the alloy to precipitate the 7' phase to a minimum of about 25 volume percent in a stable homogenous dis tribution having an effective interparticle spacing not exceeding about 5 microns at a temperature sufficiently high to prevent substantial sigma phase precipitation;

warm-working the alloy to effect an area reduction of at least percent while maintaining essentially the same volume percent and distribution of the y phase established in the aging process, providing a microdefect array of regular geometry; and

stabilizing the microdefect array of the alloy and providing additional strengthening in a final heat treatment at a temperature not exceeding the prior aging and working temperature by precipitation of an additional quantity of the y phase and by precipitation of any sigma phase as equiaxed particles.

7. The method according to claim 6 wherein:

the alloy is worked about the aging temperature.

8. The method according to claim 6 wherein:

the alloy is worked to effect an area reduction of l560 percent.

9. The method according to claim 6 wherein:

the alloy is aged at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation upon warm-working of the regular microdefect array comprising a polygonal subcell structure.

10. The method of strengthening the nickel-base superalloys of the 'y-'y type having a quantity of the 7/ phase exceeding about 25 volume percent which comprises the steps of:

heat treating the alloy to solution at least the major portion of the 'y' precipitate;

aging of the alloy at a temperature between the solvus temperature of the 7/ phase and the minimum recovery temperature to reprecipitate at least 25 volume percent of the 7 phase, based on the overall alloy composition, in a stable uniform distribution at an effective interparticle spacing not exceeding about 5 microns; working the aged alloy at about the aging temperature to effect a deformation corresponding to at least a 15 percent area reduction while maintaining essentially the same volume percent and distribution of the 7' phase established in the aging process, providing a regular array of microcrystalline imperfections; and heat treating the alloy at a temperature not exceeding the temperature of aging and working to cause precipitation of that portion of the y phase remaining in solution after aging and to cause, in those alloys containing more than 0.05 weight percent carbon, precipitation of intragranular carbides, to provide a thermally and mechanically stable array of microcrystalling imperfections.

11. The method according to claim 10 wherein:

aging is performed at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation upon working of a polygonal metallurgical substructure having the dis ocations aligned at subcell bounthe 'y'y' type enriched in those elements promoting the precipitation of topologically close-packed phases which comprises the steps of:

intermetallic heat treating the alloy to solution at least the major proportion of the intermetallic precipitates;

aging the alloy at a temperature between the solvus temperature of the 7 phase and the minimum recovery temperature, and above the precipitation temperature of the topologically close-packed phases, to reprecipitate at least 25 volume percent of the 'y phase, based on the overall alloy composition, in a stable uniform distribution at an effective interparticle spacing not exceeding about 5 microns;

working the alloy at about the aging temperature to effect a deformation corresponding to at least a 15 percent area reduction while maintaining essentially the same volume percent and distribution of the 'y' phase established in the aging process; and 4 heat treating the alloy at a temperature not exceeding the temperature of aging and working to cause precipitation of that portion of the 7 phase remaining in solution after aging, precipitation of the topologically close-packed phases as equiaxed particles and, in those alloys containing more than 0.05 weight percent carbon, precipitation of intragranular carbides, to provide a thermally and mechanically stable array of microcrystalline imperfections.

14. The method according to claim 13 wherein:

aging is performed at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation of a polygonal metallurgical substructure having the dislocations aligned at subcell boundaries to provide a regular array of defects along these boundaries.

15. The method according to claim 13 wherein;

aging is performed at a temperature sufficiently close to the minimum recovery temperature for the alloy to promote formation of warm-worked metallurgical substructure comprising a randomly nonoriented homogenous dislocation distribution.

16. The method according to claim 13 wherein:

the final heat treatment includes a sequential heat treatment, one heat treatment being selected to precipitate the topologically close-packed phases and another being selected to precipitate the 7' phase. 

2. The method according to claim 1 wherein: the initial heat treatment includes aging at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation upon warm-working of the regular microdefect array comprising a polygonal subcell structure.
 3. The method according to claim 1 wherein: the initial heat treatment includes aging at a temperature sufficiEntly close to the minimum recovery temperature of the alloy to promote formation upon warm-working of the microdefect array comprising a warm-worked metallurgical substructure consisting of a randomly nonoriented homogenous dislocation distribution.
 4. The method according to claim 2 wherein: the aging in the initial heat treatment is also above the precipitation temperature for the topologically close-packed phases.
 5. The method according to claim 3 wherein: the aging in the initial heat treatment is also above the precipitation temperature for the topologically close-packed phases.
 6. The method of processing the nickel-base superalloys of the gamma - gamma '' type having a quantity of the gamma '' phase at room temperature exceeding about 25 volume percent and, optionally, a sigma phase precipitate which comprises the steps of: heat treating the alloy to solution the precipitate; aging the alloy to precipitate the gamma '' phase to a minimum of about 25 volume percent in a stable homogenous distribution having an effective interparticle spacing not exceeding about 5 microns at a temperature sufficiently high to prevent substantial sigma phase precipitation; warm-working the alloy to effect an area reduction of at least 15 percent while maintaining essentially the same volume percent and distribution of the gamma '' phase established in the aging process, providing a microdefect array of regular geometry; and stabilizing the microdefect array of the alloy and providing additional strengthening in a final heat treatment at a temperature not exceeding the prior aging and working temperature by precipitation of an additional quantity of the gamma '' phase and by precipitation of any sigma phase as equiaxed particles.
 7. The method according to claim 6 wherein: the alloy is worked about the aging temperature.
 8. The method according to claim 6 wherein: the alloy is worked to effect an area reduction of 15-60 percent.
 9. The method according to claim 6 wherein: the alloy is aged at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation upon warm-working of the regular microdefect array comprising a polygonal subcell structure.
 10. The method of strengthening the nickel-base superalloys of the gamma - gamma '' type having a quantity of the gamma '' phase exceeding about 25 volume percent which comprises the steps of: heat treating the alloy to solution at least the major portion of the gamma '' precipitate; aging of the alloy at a temperature between the solvus temperature of the gamma '' phase and the minimum recovery temperature to reprecipitate at least 25 volume percent of the gamma '' phase, based on the overall alloy composition, in a stable uniform distribution at an effective interparticle spacing not exceeding about 5 microns; working the aged alloy at about the aging temperature to effect a deformation corresponding to at least a 15 percent area reduction while maintaining essentially the same volume percent and distribution of the gamma '' phase established in the aging process, providing a regular array of microcrystalline imperfections; and heat treating the alloy at a temperature not exceeding the temperature of aging and working to cause precipitation of that portion of the gamma '' phase remaining in solution after aging and to cause, in those alloys containing more than 0.05 weight percent carbon, precipitation of intragranular carbides, to provide a thermally and mechanically stable array of microcrystalling imperfections.
 11. The method according to claim 10 wherein: aging is performed at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation upon working of a polygonal metallurgical substructure having the dislocations aligned at subcell boundaries to provide a regular array of defects along these boundaries.
 12. The method according to claim 10 wherein: aging is performed at a temperature sufficiently close to the minimum recovery temperature for the alloy to promote formation upon working of a regular array of microcrystalline imperfections comprising a warm-worked metallurgical substructure comprising a randomly nonoriented homogenous dislocation distribution.
 13. The method of processing the nickel-base superalloys of the gamma - gamma '' type enriched in those elements promoting the precipitation of topologically close-packed intermetallic phases which comprises the steps of: heat treating the alloy to solution at least the major proportion of the intermetallic precipitates; aging the alloy at a temperature between the solvus temperature of the gamma '' phase and the minimum recovery temperature, and above the precipitation temperature of the topologically close-packed phases, to reprecipitate at least 25 volume percent of the gamma '' phase, based on the overall alloy composition, in a stable uniform distribution at an effective interparticle spacing not exceeding about 5 microns; working the alloy at about the aging temperature to effect a deformation corresponding to at least a 15 percent area reduction while maintaining essentially the same volume percent and distribution of the gamma '' phase established in the aging process; and heat treating the alloy at a temperature not exceeding the temperature of aging and working to cause precipitation of that portion of the gamma '' phase remaining in solution after aging, precipitation of the topologically close-packed phases as equiaxed particles and, in those alloys containing more than 0.05 weight percent carbon, precipitation of intragranular carbides, to provide a thermally and mechanically stable array of microcrystalline imperfections.
 14. The method according to claim 13 wherein: aging is performed at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation of a polygonal metallurgical substructure having the dislocations aligned at subcell boundaries to provide a regular array of defects along these boundaries.
 15. The method according to claim 13 wherein; aging is performed at a temperature sufficiently close to the minimum recovery temperature for the alloy to promote formation of warm-worked metallurgical substructure comprising a randomly nonoriented homogenous dislocation distribution.
 16. The method according to claim 13 wherein: the final heat treatment includes a sequential heat treatment, one heat treatment being selected to precipitate the topologically close-packed phases and another being selected to precipitate the gamma '' phase. 